1. Introduction

Mid-infrared (MIR) light sources, which operate in the range of wavelengths spanning from 3 μm to 25 μm, present extensive potential applications in the fields of medicine, environment pollution monitoring, biology, agriculture, and security [1, 2] because the resonant oscillation frequencies of numerous molecular bonds fall within MIR wavelength range. For instance, the absorption intensity of the basic vibrations of the molecules in the MIR region is three to five orders of magnitude higher than that in the near-infrared region [3], and the sensitivity of the gas sensor can be markedly improved by the direct measurement of the gas in the MIR band. With advances in infrared fiber and detector technology, the development of low-cost, high-brightness MIR light sources covering the 3–5 μm band has become the key to developing MIR fiber sensing technology. Ari et al. studied the mid-IR luminescence properties of Pr³⁺ ion-doped GaGeSbSe glass and fiber [4]. The mid-IR luminescence covers the CO₂ absorption band at 4.3 μm and can be used in an environmental monitoring sensor for CO₂ underground storage. Starecki et al. reported fluorescence absorbance-based mid-IR optical sensor for CO₂ detection [5]. Schweizer found that the use of Pr³⁺ ions in 3.4 μm fluorescence detected of butane and natural gas concentration, by increasing the length of the chamber and the detection of the edge of the absorption band, the minimum detection concentration of up to 50 ppm [6]. However, the development of the rare earth (RE)-doped chalcogenide glass fiber is limited by low solubility and low emission efficiency. Meanwhile, impurities in the chalcogenide glass matrix, which are mainly due to hydrogen and oxygen, not only deteriorate glass quality, increase the tendency for crystallization, and causing selective impurity absorption and scattering, but also interact with RE-ions, thereby inducing ion–ion relaxation and quenching of the photoluminescence [7, 8].

In accordance with recent investigations of chalcogenide glasses for active fiber optics, elements Ga, In, and I improve the solubility of RE in different glass matrices [9–11]. Seddon et al. found that the introduction of Ga in a Ge–As–Se glass system yields better dispersibility than In, which is beneficial to improve the Pr³⁺ ion doping concentration [11]. Shiryaev et al. reported that Ge–Sb–Se–In–I glass can improve the solubility of rare earth (RE) elements compared with other sulfur groups owing to the molecular structure of the added In and I [12]. Crystalline glass ceramics can further improve the photoluminescence of glass [13]. Wang et al. studied the EELS spectrum and found that the local refractive index of Dy³⁺ ions was enhanced by heat treatment and the precipitation of Ga₂S₃ nanocrystals in glass [14].

In this work, the Ge–Ga–Sb–Se–I glass was selected as the matrix. The luminescence properties were studied with the increase in GaI or Pr³⁺ ion concentration. Se purification was also attempted to diminish the effect of contamination. Raman spectroscopy was used to characterize the internal structure of the glass. Pr³⁺-doped chalcogenide glass ceramics were prepared by heat treatment, and the emission intensity at 4.75 μm increased significantly with prolonged heat treatment duration.

2. Experimental

Ge₂₂Sb₈Se_70−2x(GaI)_x–0.2Pr (x = 1, 2, 3, 4, and 5, in mol%) and Ge₂₂Sb₈Se₆₂(GaI)₄–yPr (y = 0.1, 0.2, 0.3, 0.5, 1.0, and 1.5, in mol%) glasses were prepared using the standard melt-quenching method. GaI₃ (6N) was used as starting material to prevent the sublimation of element I₂. Ge (5N), Sb (5N), Se (5N) and Pr (3N) were used as starting materials. Raw materials (20 g) were weighted in a glove box and loaded into a low -OH, high-purity fused silica tube with an inner diameter of 12 mm. After evacuation to 10⁻⁷ torr, the tube was heated at a temperature of 100 °C for 1 h to remove the surface water. The tube was sealed by using a gas torch, placed in a rocking furnace, and heated at 950 °C for 10 h at minimum. Finally, the tube containing the melt was quenched in water, and the formed glass was annealed below glass transition temperature (T_g–20 °C) for 5 h. Glass samples were cut into disks (~2 mm) and polished to high optical quality for further tests.

Glass transition temperature (T_g) was measured at a rate of 10 K/min by using a differential scanning calorimeter (Q2000, TA) under N₂ gas protection. The refractive index (n) at 1.7–17 μm was determined using an infrared variable angle spectroscopic ellipsometer (IR-VASE Mark II, J. A. Wollam). The optical absorption spectra were obtained with a spectrophotometer (Lambda 950, Perkin-Elmer) in the range of 400–2500 nm and an infrared spectrometer (FT-IR, Nicolet 380) in the range of 2.5–25 μm, respectively. X-ray diffraction patterns (XRD) were recorded under ambient conditions by using a Bruker D2 X-ray diffractometer. MIR fluorescence spectra in the range of 3–5.5 μm were recorded by an InSb detector combined with an FLS980 fluorescence spectrometer at room temperature. A solid-state laser at 1.5 μm was used as excitation source. All test conditions, including pump power and sample position, were identical to ensure that the measured fluorescence intensities are comparable. The morphology of crystals inside the glass ceramics were observed by SEM and TEM. Crystal formation was calculated from HRTEM images. All measurements were performed at room temperature.

3. Results and discussion

The measured T_g, density (ρ), and refractive index at 4.5 μm of the glasses are listed in Table 1. With the addition of GaI, the T_g gradually increased due to the increase in tetracoordinate Ga in the glass matrix. As the GaI concentration increased, the n decreased slightly. This response is relative to the ionic polarization (p) of constituent elements and the density of the glass sample, but the relationship between the two factors is complicated [15]. Given the similar glass composition, the same ion polarizability can be considered. The density of glass exerts a considerable impact on the n. The n decreased as the density also decreased (Table 1).

Table 1. Glass compositions with corresponding T_g, density, and refractive index at 4.5 μm

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3.1 Absorption spectra

Figure 1(a) shows the absorption spectra of Pr³⁺-doped Ge–Sb–Se–Ga–I glasses in the region of 0.5–5.5 μm. The main electronic absorption bands at 1.48, 1.63, 2.04, 2.33, and 4.5 μm corresponded to transitions from the ground state ³H₄ to the ³F₄, ³F₃, ³F₂, ³H₆, and ³H₅ upper levels, respectively. The absorption band at 2.9 μm was assigned to O-H impurity vibrational absorption. The absorption band observed at 4.5 μm also encompassed H-Se extrinsic impurity vibrational absorption. Figure 1(b) presents the electronic energy level diagram of Pr³⁺ ion and shows the MIR absorptive transitions and potential radiative transitions. Notably, (Figs. 1(a) and 1(b)) that Pr³⁺ ions exhibited absorption bands at 1.5 (³H₄→³F₄, ³F₃) and 2 μm (³H₄→³F₂, ³H₆), indicating that Pr³⁺-doped Ge–Sb–Se–Ga–I glass can be pumped by commercially available sources at ~1.5 and ~2 μm [10].

 

Fig. 1 (a) Absorption spectra of Pr³⁺ doped Ge–Sb–Se–Ga–I bulk glasses; (b) Pr³⁺ level and related transitions

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3.2 Luminescence properties

Broadband photoluminescence was observed for prepared glasses in the wavelength range of 3.5–5.5 μm under 1.55 μm laser diode (LD)excitation (Fig. 2). According to theoretical evaluation, the nature of this band is associated with several radiative transitions. Strong radiative transitions of ³F₂→³H₅, ³H₆ →³H₅ and ³H₅→³H₄ mainly contributed to the emission [16, 17]. The dips in emission spectra were due to CO₂ absorption (4.25 μm) in air and Se-H impurity absorption (4.57 μm). Moreover, the emission intensity at 3–5.5 μm increased with addition of GaI as shown in the Fig. 2(a). The maximum emission was obtained in the sample containing mostly GaI contents. The strongest fluorescence intensity was achieved at 4 mol% GaI. It is suggested that there are preferential spatial correlations between REI and the GaS₄ tetrahedral even at low Ga-doping levels [18]. Furthermore, the phenomenon that Ga complexes the RE-ions by forming [RE³⁺-S-Ga³⁺] type species has already been confirmed [19]. Therefore, according to the experience mentioned above and with similar chemical properties of S and Se, when GaI content reached 4 mol%, it can be supposed that the introduction of Ga to the GeSbSeI structure has dispersed Pr³⁺ ions in the glass. This dispersion is conducive to fluorescence.

 

Fig. 2 MIR emission spectra of Pr³⁺-doped samples under 1.55 μm nm laser excitation (a) with increase in GaI concentration and (b) with increase of Pr³⁺ ion concentration.

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Figure 2(b) shows the infrared emission in 4 mol% GaI glass sample as a function of Pr³⁺ concentration. The bandwidth of the emission spectra increased with the increase in Pr³⁺ concentration because of the enhanced interaction between Pr³⁺ excited state energy levels. With the increase in Pr³⁺ concentration, ion spacing decreased, and the transition probability of Pr³⁺ intercourse relaxation increased. Notable maximum emission intensity at 3–5.5 μm was obtained until 1.5 mol% Pr³⁺ was introduced into the glass matrix. The concentration quenching effect was not observed in the 1.5 mol% Pr³⁺-doped sample, but the glass surface became extremely rough. Figure 1(b) shows that the corresponding spectra of the (³F₂, ³H₆→³H₅ and ³H₅→³H₄) two transitions a large overlap; thus, the transition level at ³F₂ and ³H₆ can be effectively reabsorbed and increase the ³H₅ level reversal of a number of particles to promote the broadband emission in the range of 3–5.5 μm [11].

Figure 3 shows the absorption and emission spectra of the Pr³⁺-doped glass with/without Se-purification, respectively. Impurities can weaken the emission intensity by multiphonon relaxation. There are different populations of RE-ions in a selenide glass host. There are other smaller populations of RE-ions whose radiative behavior is dominated instead by the maximum phonon energies of local impurities, which can lead to fast non-radiative decay for some fraction of the RE-ion population. These other non-radiative decay processes were proposed to be energy transfer to vibrational impurities. Among all the raw materials, the Se is the most easily oxidized raw material, and oxides are the most electronegative anions. The pre-purification of Se can largely decrease the oxygen contents in the glass. For selenide glasses, the H-Se impurity absorption band overlaps with Pr³⁺ emission bands in the 4–5 μm wavelength region. Thus, only one phonon is required to de-activate the electrons of Pr³⁺-ion excite state to Se-H impurity. It is difficult to distinguish the change of Se-H after purification by absorption spectrum and fluorescence spectrum (Fig. 3). The bond of Se-H is difficult to remove completely and only part of it can be removed. Given that impurity multiphonon relaxation already caused fast non-radiative decay, the measured fluorescent intensities are lower than those with Se-purification [8]. Reducing impurities can improve the electron population of the excited state and luminous intensity of Pr³⁺. Thus, glass purification is very important for efficient MIR fluorescence.

 

Fig. 3 Pr³⁺-doped Ge₂₂Sb₈Se₆₆(GaI)₂ bulk glass and Se purification of Pr³⁺-doped Ge₂₂Sb₈Se₆₆(GaI)₂ bulk glass: (a) absorption spectra; (b) photoluminescence spectra.

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3.3 Raman spectra

Figure 4 illustrates the normalized Raman spectra of the Ge–Sb–Se–Ga–I glasses. In the Ge–Sb–Se–Ga–I glasses, the vibrational modes of structural units containing gallium were not evident. The vibrational modes were likely to overlap the structural units containing germanium owing to a slight difference between Ge and Ga atomic weights. The Raman spectra of these glasses can be interpreted on the basis of known date for the Ge–Sb–Se glass structure, which consisted mainly of corner-sharing and edge-sharing GeSe₄ tetrahedrons. The broad band in the 150–240 cm⁻¹ regions indicates the superposition of several bands. The bands at 200 and 214 cm⁻¹ were ascribed to the symmetrical stretching vibration of the Ge–Se bonds in the corner-sharing GeSe_4/2 tetrahedron and edge-sharing GeSe_4/2 bi-tetrahedron, respectively [20]. The bond that may be caused by vibrations of the Sb–Se bonds in SbSe_3/2 pyramids are linked by bridging atoms of Se at 195 cm⁻¹ [21]. With the addition of GaI, Se atoms were gradually consumed. A shoulder peak appears at 176 cm⁻¹, which was assigned to vibrations of Ge(Ga)Se_6/2 structural units with Ge(Ga)–Ge(Ga) bonds [22]. The band in the 235–320 cm⁻¹ range with low intensity consisted of a minimum of five overlapping bands: at 250 cm⁻¹ due to Se₈ rings [22]; at 256 cm⁻¹ due to the vibrations of Se_n; at 266 cm⁻¹ due to Se–Se chains [20, 23]; and at 275 and 307 cm⁻¹ due to GeSe₂. The bands in the 235-280 cm⁻¹ range presented a noticeted a noticeable decrease with addition of GaI. This was ascribed to the gradual depletion of Se elements. Meanwhile, the band in the 280-325 cm⁻¹ range which resulted from F₂ asymmetric vibration of [GeSe₄] and/or [GaSe₄] tetrahedra increased with addition of GaI. The incorporation of gallium as modifiers into chalcogenide glassy networks could dramatically increase the solubility of RE ions due to the presence of edge-sharing [GaSe₄] tetrahedral structure. Bonds between like atoms were found in both nonstopchiometric regions (Se rich). This disorder of the glasses is accompanied by presence of ‘wrong’ bonds (Se-Se bonds) [24]. As the glass gets closer to stoichiometry, the defects inside the glass will be reduced. Closer to the stoichiometric composition is more conducive to uniform dispersion of Pr³⁺ in the glass matrix. The defects absorbing energy will be reduced. and the fluorescence will be increased. The glasses exhibit low phonon energy which could help reduce the multiphonon relaxation of excited states. Therefore, the MIR luminescence efficiency of RE ions in these glasses can be improved.

 

Fig. 4 Normalized Raman spectra of Ge₂₂Sb₈Se_70−2x(GaI)_x(x = 1, 2,3,4,5)–0.2Pr glass.

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3.4 Crystallization behavior

Crystallization can further enhance emission intensity. Ge₂₂Sb₈Se₆₂(GaI)₄–0.5Pr glass was selected to study the influence of crystallization on the emission of Pr³⁺ ion. From the DSC curve, the characteristic temperatures of T_g = 282 °C, T_x = 400 °C, and T_p = 423 °C were determined. Based on the DSC results and previous studies [25–27], a relatively low T_HT of 312 °C (T_g + 30 °C) was selected for heat treatment. The base glasses were placed in a muffle furnace for different durations ranging from 5 h to 48 h.

Figure 5(a) shows the XRD patterns of samples heat treated for 5 h to 48 h. The appearance of crystallization was observed for approximately 10 h, and the crystal peaks were assigned to GeSe₂ (JCPDF 30-595) crystals. Meanwhile, the Sb₂Se₃ crystals (JCPDF 15-861) appeared at 24 h. Both GeSe₂ and Sb₂Se₃ crystals precipitated at 48 h. The SEM images shown in Fig. 5(b) indicate that the base glass is not crystallized, presenting a smooth surface. Figure 5(c) indicates that the glass was crystallized completely at 312 °C for 24 h. A TEM image of Ge₂₂Sb₈Se₆₂(GaI)₄–0.5Pr heat treated for 24 h is shown in Fig. 5(d). By carefully measuring the interplanar distance, two types of crystals can be distinguished. The lattice plane distances are 0.3567 and 0.3837 nm, which correspond to the d₀₂₃ and d₀₁₃ of GeSe₂ (JCPDF 30-595), respectively. GeSe₂ crystals have many crystal faces, and they are just different crystal faces of the same crystal.

 

Fig. 5 (a) XRD patterns of Ge₂₂Sb₈Se₆₂(GaI)₄–0.5Pr base glass and glass ceramics obtained by heat treatments at 312 °C for different durations, respectively, (b) SEM images for base sample at × 50000, (c) SEM images for sample at 312 °C for 24 h at × 50000, and (d) TEM image of the Ge₂₂Sb₈Se₆₂(GaI)₄–0.5Pr sample crystallized at 312 °C 24 h.

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The MIR emission spectra of Pr³⁺-doped glass and glass ceramic samples obtained by heat treatments were collected with a laser excitation at 1.55 μm. As shown in Fig. 6, the emission bands which correspond to ³F₂→³H₅, ³H₆→³H₅ and ³H₅→³H₄ transitions were observed in all samples. With the increase in heat treatment durations, fluorescence intensity gradually increased. The maximum fluorescence intensity, which is two times higher than that of the base glass, was reached at 10 h. Fluorescence intensity was gradually reduced by prolonging the heat treatment durations. The Pr³⁺ ions in transparent Ge₂₂Sb₈Se₆₂(GaI)₄ were surrounded by an ionic environment with low phonon energy, which facilitated the MIR emission because nonradiative transition is proportional to phonon energy [28]. This behavior was further confirmed by the emission peaks, which became sharper in the Ge₂₂Sb₈Se₆₂(GaI)₄ sample than those in the precursor glass. The evolution of MIR luminescence with heat treatment durations were mainly influenced by crystallization. At the initial stage, small crystals were precipitated into the glass matrix, and the crystalline environment promoted the emission. With the increasing heat treatment durations, the size of the crystals increased. Light scattering by crystals diminishes emission.

 

Fig. 6 Infrared emission spectra of Ge₂₂Sb₈Se₆₂(GaI)₄–0.5Pr base glass and glass ceramics obtained by heat treatments at 312 °C for different durations under 1550 nm laser excitation, respectively.

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Conclusions

Ge₂₂Sb₈Se_70−2x(GaI)_x–0.2Pr (x = 1, 2, 3, 4, 5) and Ge₂₂Sb₈Se₆₂(GaI)₄–yPr (y = 0.1, 0.2, 0.3, 0.5, 1.0, 1.5) glasses were synthesized by conventional melt-quenching technique. MIR (3–5 μm) fluorescence spectra of glasses were observed and compared under excitation of 1550 nm laser. The highest emission intensity was obtained in Ge₂₂Sb₈Se₆₂(GaI)₄–0.2Pr glass. Raman spectroscopy showed that the Pr³⁺-doped Ge–Sb–Se–Ga–I glasses mainly consisted of GeSe_4/2 tetrahedron, Se₈^‒ rings, Se–Se chains, and SbSe_3/2^‒ pyramids structural units. IR transparent chalcohalide glass ceramics with composition of Ge₂₂Sb₈Se₆₂(GaI)₄ doped with 0.5 mol% Pr³⁺ ions were prepared by heat treating the base glass at 312 °C for different durations. XRD and TEM results indicated the appearance of GeSe₂ and Sb₂Se₃ nanocrystals in the samples. MIR emission showed significant increase in luminescence of 3 μm to 5.5 μm in the sample relative to that of the precursor glass. The glass treated at 312 °C for 10 h reached the maximum emission intensity. Luminescence characteristics showed that the Pr³⁺-doped Ge–Sb–Se–Ga–I glasses exhibit high potential for the development of MIR fiber lasers and amplifiers.